Hard weld overlays resistant to re-heat cracking

ABSTRACT

A hard weld overlay which is resistant to cracking when re-heated, and a method for designing such alloys, is disclosed. The alloys are able to resist re-heat cracking through prevention of the precipitation and/or growth of embrittling carbide, borides, or borocarbides at elevated temperatures. In one embodiment, the thermodynamics of the alloy system possess only primary carbides and secondary ferrite carbides.

INCORPORATION BY REFERENCE TO ANY PRIORITY APPLICATIONS

Any and all applications for which a foreign or domestic priority claimis identified in the Application Data Sheet as filed with the presentapplication are hereby incorporated by reference under 37 CFR 1.57.

BACKGROUND

1. Field

This invention relates to hard coatings and weld overlays used toprotect surfaces from wear.

2. Description of the Related Art

The hardfacing process is a technique used to protect a surface fromwear. Typical methods of hardfacing include the various methods ofwelding, GMAW, GTAW, PTA, laser cladding, submerged arc welding, openarc welding, thermal spray, and explosive welding. In certainapplications, it is desirable for the hardfacing coating to be free ofcracks. Hardbanding, the process of applying a hardfacing layer to theouter diameter of tool joints on a drill string, is an example of anapplication where cracks are undesirable. Cracks can allow forcorrosion, create welding difficulties when re-building the hardbandinglayer, and allow for the propagation of cracks from the hardfacing layerinto the substrate material resulting in the failure of the drill pipeitself. Preventing cracking can be achieved in hardbanding materials byincreasing the toughness of the hardfacing alloy used. However, hardnessand toughness are inversely related material properties. Thus, in orderto prevent cracking the hardness is sacrificed. Typical non-crackinghardfacing materials deposited via the GMAW process for the purposes ofhardbanding possess hardness in the range of 50-60 HRC. Crackinghardfacing materials such as chromium carbide can exhibit hardnessessignificantly above 60 HRC, in the range of 61-69 HRC.

Several modes of cracking are known to occur in hardbanding. Three typesof cracking occur during welding, or slightly after (1 s-180 s) thewelding as been completed. Cross checking is defined as a large crackwhich spans across the entire weld bead width, and can occur during thedeposition of a single bead. The two other forms of cracking, dipcracking and circumferential cracking are associated with the re-heatingof an existing bead. Dip cracking occurs during the welding of a singlebead.

During the hardbanding process, a 1″ wide weld bead is deposited onto arotating tool joint such that it covers the entire circumference of thejoint when completed. The weld is completed when joint has made one fullrevolution during the weld process, such that new weld material isdeposited directly on top of existing weld material. This overlap causesthe existing weld material to re-heat. The re-heating of the existingweld material can cause dip cracking in the existing bead 0-2 inchesaway from the overlap zone.

Circumferential cracking occurs when multiple bands are welded next toeach other, as is customary in the hardbanding process and otherhardfacing processes. In the hardbanding process, it is customary tooverlap one bead with subsequent weld passes by ⅛″ to ¼″. This slightoverlap between neighboring beads re-heats the existing bead and canlead to circumferential cracking.

The thermodynamic profile of a standard hardbanding alloy,Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54)[alloy 1], is displayed in FIG. 1. Alloy 1 shows the solidification ofthe liquid (phase 9) to austenite (phase 7) before transforming toferrite (phase 1), in addition to the solidification of several primarycarbides (phase 2, 3, 4, 6, and 8) However, an important part of thethermodynamic behavior of alloy 1 is thermodynamic stability of VC(phase 4). VC is thermodynamically stable from room temperature toslightly below 1200 K, which categorizes this phase as a secondaryaustenite carbide. Upon initial solidification of the weld, the primarycarbides and the iron matrix solidify. However, due to the rapid coolingin a weld bead, the VC phase is unlikely to form due to sluggishkinetics at these lower temperatures. The inability for VC toprecipitate and grow to a stable size upon the initial deposition of theweld bead produces a super saturation of carbon in the iron-basedmatrix. As the weld alloy is re-heated during the continuation of theweld process, the temperature in the re-heat zone enters into the VCstable temperature range and VC can form from the ferrite or austenite.The newly formed VC introduces additional stress in the alloy and causescracking when the weld rapidly cools again.

SUMMARY

A hard weld overlay which is resistant to cracking when re-heated, and amethod for designing such alloys, is disclosed. The alloys are able toresist re-heat cracking through prevention of the precipitation and/orgrowth of embrittling carbide, borides, or borocarbides at elevatedtemperatures. Reheating of existing hard weld overlays exposes thematerial to elevated temperatures and occurs frequently in hard claddingoperations due to overlapping and multiple pass overlays. In most hardweld overlay processes, a certain fraction of carbon and boron istrapped in the matrix due to slow cooling. Upon re-heating, the trappedcarbon and boron is released to form carbides and borides. Theprecipitation and growth of the carbides in combination with the graingrowth in the steel matrix is known to cause stress in themicrostructure and lead to cracking. By controlling the thermodynamicsof the boride and carbide phases, it is possible to create an alloywhich is less prone to the growth of new carbides and borides duringre-heating, and is thus less prone to cracking. When designing suchalloys, different carbides and borides can be classified into threedistinct groups: primary carbides, secondary austenite carbides, andsecondary ferrite carbides.

Primary carbides are thermodynamically stable at temperatures higherthan or within 50° C. of the initial solidification temperature of thematrix. Secondary austenite carbides become thermodynamically stable attemperatures above the ferrite to austenite transition temperature butmore than 50° C. below the initial solidification temperature of thematrix. Finally, secondary ferrite carbides are only thermodynamicallystable at temperatures below the austenite to ferrite transition.

In one embodiment, the thermodynamics of the alloy system possess onlyprimary carbides and secondary ferrite carbides. In such as system, theprimary carbides will solidify prior to the solidification of theaustenite, and the remaining carbon and boron will likely be trapped inthe ferrite as the alloy eventually cools to room temperature, i.e. thesecondary ferrite carbides are kinetically unable to precipitate andgrow to significant levels. As the weld alloy is re-heated during thecontinuation of the weld process, the temperature in the re-heat zoneenters into the austenite region for the steel. At these temperatures,the secondary ferrite carbides are not thermodynamically stable and donot form. Grain growth of the matrix occurs, but no new significantcarbide formation has occurred. As the alloy rapidly cools again, theferrite secondary carbides are unable to form due to sluggish kineticsand the majority of the carbon and boron is once again trapped in thematrix.

In a preferred embodiment, the primary carbides are at least one of:titanium boride (TiB₂) or Niobium carbide (NbC).

Thermo-Calc is a powerful software package used to perform thermodynamicand phase diagram calculations for multi-component systems of practicalimportance. Calculations using Thermo-Calc are based on thermodynamicdatabases, which are produced by expert evaluation of experimental datausing the CALPHAD method.

TCFE7 is a thermodynamic database for different kinds of steels,Fe-based alloys (stainless steels, high-speed steels, tool steels, HSLAsteels, cast iron, corrosion-resistant high strength steels and more)and cemented carbides for use with the Thermo-Calc, DICTRA and TCPRISMAsoftware packages. TCFE7 includes elements such as Ar, Al, B, C, Ca, Co,Cr, Cu, H, Mg, Mn, Mo, N, Nb, Ni, 0, P, S, Si, Ta, Ti, V, W, Zr and Fe.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1:

Phase evolution ofFe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54)[alloy 1].

FIG. 2:

Phase evolution ofFe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54)[alloy2].

FIG. 3:

Phase evolution diagram ofFe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂[alloy 3].

FIG. 4: Elemental concentration in NbC phase.

FIG. 5: FCC to BCC transition temperature in selected hardbanding alloys

FIG. 6:

Phase evolution diagram ofFe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.5)Si_(0.59)Ti₁V_(0.54).

FIGS. 7A-B: Optical microstructures at 500× of alloy 5 (7A) and alloy 6(7B).

DETAILED DESCRIPTION

In one embodiment, the thermodynamic properties of the alloy arecalculated using the CALPHAD method. A preferred embodiment uses theThermo-Calc software to perform these calculations.

Non-Cracking Trait 1:

In one embodiment, all of the carbide, boride, and boro-carbide phasesare primary carbides. Thus, they are thermodynamically stable at therelatively high temperatures as defined previously. An alloy whichpossesses this thermodynamic profile is more resistant to cracking thanconventional hardfacing materials. As an alloy of this type is initialdeposited in the form of a weld bead, the primary carbides begin toprecipitate and grow during the initial solidification of the material.Typically, a large fraction of primary carbides precipitate prior to thesolidification of the matrix. This solidification is advantageous forimproving crack resistance, in that the existing primary carbides do notinflict high stresses on solidifying austenite or during thetransformation of austenite to ferrite. The formation of primarycarbides effectively reduces the total carbon in the solidifyingaustenite such that is less likely for the iron-based matrix to becomesuper saturated with carbon. This aids in final structure of the metalbeing ferritic as opposed to austenitic, and aids in the resistance ofcracking during re-heating.

In conventional hardfacing materials, the iron-based matrix is oftensuper saturated with carbon. Upon re-heating, the carbon is allowed todiffuse throughout the microstructure and form carbides. As the matrixtransforms to austenite and the grain size increases, these newly formcarbides cause stresses on the microstructure of the material, which canlead to cracking in the hardfacing material.

In the described embodiment where all carbides, borides, andboro-carbide phases are primary carbides this described crackingmechanism is avoided. Upon re-heating in the described embodiment, thematrix does not form new carbides and thus stresses are avoided as thematrix transforms and grows. An example alloy [alloy 2],Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54),that demonstrates this phenomenon is shown in FIG. 2. This diagram showsthe solidification of the liquid (phase 7) into austenite (phase 6),which ultimately transforms to ferrite (phase 2). This is the commonfeature of the equilibrium solidification pathway for most steels. Theunique components of this alloy are the solidification of the borides,carbides, and borocarbides (phases 1, 3, 4, and 5). All of these phasescan be defined as primary carbides as they form at high temperaturesclose to the solidification temperature of the austenite phase.

In this preferred embodiment, the primary carbides are TiB2 (phase 1),Cr2B (phase 3), NbC (phase 4), and (Fe,Cr)3B2 (phase 5).

In one embodiment the reheat temperature range is 800° C. to 1300° C.

In a preferred embodiment the reheat temperature range is 900° C. to1200° C.

In a still preferred embodiment the reheat temperature range is 1000° C.to 1100° C.

Non-Cracking Trait 2:

In another embodiment, the mole fraction of all the carbide phasesremain thermodynamically stable within the temperature range defined asthe re-heat zone. In a preferred embodiment, stability is defined as amole fraction which does not vary by more than 25%; in a still preferredembodiment stability is defined as a mole fraction which does not varyby more than 10%, in a still preferred embodiment, stability is definesas a mole fraction does not vary be more than 5%.

Carbides which are thermodynamically stable within the re-heat zone arebeneficial for the purposes of creating an alloy which is resistant tore-heat cracking. In the case of a cracking prone alloy, the re-heatingof the alloy can cause the precipitation and/or growth of additionalcarbide or the dissolution and shrinking of existing carbides. Growingor re-precipitation of carbides causes stresses in the matrix asdescribed previously. The dissolution of carbides can also bedetrimental as it increases the carbon and/or boron in the iron-basedmatrix. This increase in carbon in the matrix can cause other carbidesto precipitate or grow causing stresses in different regions of themicrostructure, or it can lead to supersaturation of carbon in thematrix which can make the material prone to re-heat cracking.

FIG. 2 depicts the thermodynamics of an alloy which possess the carbideswhich have a mole fraction that is thermodynamically stable within thereheat zone. As shown, there are no phase transformations or large phasemole fraction variations within the reheat zone. The primary carbidephases (1, 3, 4, and 5) are all stable from the austenite solidificationtemperature to temperatures below the reheat zone. When an alloy of thisphase structure is re-heated, the carbides are stable and do not grow ordissolve. This prevents additional stress in the weld and cracking canbe avoided.

In another embodiment, all of the secondary carbides are onlythermodynamically stable below the reheat zone.

Non-Cracking Trait 3:

An alloy which possesses the thermodynamics of this embodiment isresistant to cracking in the re-heat zone. The solidification routine ofsuch an alloy when initially deposited is similar to previouslydescribed: the Fe-based matrix and primary carbides solidify to form themicrostructure. The secondary carbides are kinetically unable to formdue to the rapid cooling of the process, leaving the Fe-based matrixsupersaturated with carbon and/or boron. However, as the temperature ofthe material is increased into the reheat zone, the secondary carbidephase is not thermodynamically stable so it does not form. The materialthen cool rapidly down to room temperature, and the secondary carbidephase is once again unable to precipitate due to sluggish kinetics.

A preferred embodiment, AlloyFe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂, isshown in FIG. 3. As shown, Phase 8, is a secondary carbide phase whichis only thermodynamically stable below the reheat zone. Phase 8 isunlikely to form during the original deposition of the weld bead, andunlikely to form as the material is reheated. This embodiment allows thealloy to be supersaturated with carbon, increasing hardness, but stillmaintains crack resistance.

Non-Cracking Trait 4:

In another embodiment, A selection of the carbides don't contain morethan 50% Fe. During reheating in the weld bead, the Fe-rich carbides canform much easily than other carbide. This phenomenon occurs because thematrix is Fe-rich and carbon has a much higher likelihood of diffusinginto a region of the microstructure where Fe is free to react andprecipitate new carbides. Furthermore, as the newly precipitatedcarbides or existing carbides are driven to grow in the alloy, theability to utilize the large availability of Fe as opposed to lowerconcentration alloying elements will increase the growth rate of suchcarbides. Carbides which are more likely to precipitate and capable ofgrowing rapidly in the re-heated alloy will make the alloy moresusceptible to re-heat cracking.

FIG. 4 shows the variation of the mole fraction of each element in NbC,which is a common carbide in the presented hardfacing alloys. The NbCphase contains primarily Nb and C with a slight amount of V, but traceconcentrations of Fe. Such a carbide will be unlikely to grow any largerduring the reheating of the weld, because both Nb and V will berelatively scarce around the local region of the carbide.

In one preferred embodiment, all of the secondary carbide phases don'tcontain more than 50% Fe.

In a second preferred embodiment, all of the primary carbide phasesdon't contain more than 50% Fe.

In a still preferred embodiment, the carbide phases precipitating in thealloy consist of at least one of TiB₂, CrB₂, NbC, WC, MoB₂, and/or VC.

Non-Cracking Trait 5:

In another embodiment, the alloy is designed such that the fccaustenite/bcc ferrite transition temperature is not within the RZ.Avoiding this significant phase transformation at the RZ can minimizethe stress in the microstructure and make the alloy less prone to reheatcracking. By avoiding the FCC to BCC transition upon re-heating, thealloy will be more capable of handling the stresses created by newlyprecipitated carbides or growth of existing carbides. FIG. 5demonstrates how the transition temperature of the hardfacing alloy canbe controlled by compositional variation.

In another embodiment, the RZ is shifted by adjusting the weldingparameters used in the weld process in order to avoid the fccaustenite/bcc ferrite transition temperature in a particular alloy. Thefcc austenite/bcc ferrite transition is the biggest phase transformationin the steel and can introduce significant stress causing cracking.

FIG. 5 shows the relationship between the fcc austenite/bcc ferritetransition temperature vs. carbon content. We can know what kind ofmicrostructure (ferrite, austenite or martensite) will occur afterwelding by calculating the fcc austenite/bcc ferrite transitiontemperature. We can also adjust the fcc austenite/bcc ferrite transitiontemperature by changing some elements, then obtain the optimummicrostructure.

Non-Cracking Trait 6:

In another embodiment, carbides do not form in the austenitic zone ofthe alloy during re-heating. Carbides which become stable in theaustenitic zone can precipitate and/or grow upon reheating of the alloywhen the matrix is austenitic. When the alloy is in the austenite phasegrain growth is typical and carbides typically precipitate along theprevious grain boundaries of the initially deposited ferrite matrix.Therefore, the carbides which have precipitated in the austenite are nowlocated in the center regions of the matrix grains. As the alloy coolsand transforms back to ferrite, the newly grown carbides in the centerof the grains can cause stress on the microstructure and create cracks.An alloy which avoids the precipitation of carbides in the austenitezone is shown in FIG. 6. The VC, phase 3, is not thermodynamicallystable in the austenite region (phase 6). Thus, any precipitation of VCdo to the re-heating of the weld occurs after the alloy has transitionedfrom BCC to FCC upon heating and back to BCC upon cooling. Therefore,the newly formed carbide is not present during the potentiallystress-inducing, and thereby crack prone, solid state transition.

In a one embodiment, the hardfacing alloy is Fe-based containing one ormore of the following alloying elements B, C, Cr, Mn, Mo, Nb, Si, Ti, W,and V with additional impurities known to be present due tomanufacturing procedures and possesses one of the preferred non-crackingtraits described in this disclosure.

In a preferred embodiment, this hardfacing alloy is in the form of acored welding wire.

In another preferred embodiment the hardfacing alloy composition, asdefined by the composition of the feedstock material or the depositedcoating, is given in weight percent by the following range:Fe_(ba1)C_(0.5-4)B₀₋₃Mn₀₋₁₀Al₀₋₅Si₀₋₅Ni₀₋₅Cr₀₋₃₀Mo₀₋₁₀V₀₋₁₀W₀₋₁₅Ti₀₋₁₀Nb₀₋₁₀

In a still preferred embodiment the hardfacing alloy composition, asdefined by the composition of the feedstock material or the depositedcoating, is given in weight percent by the following range:Fe_(ba1)C₁₋₂B_(1-2.5)Mn₁₋₂Al_(0-0.5)Si_(0-1.5)Ni_(0-0.2)Cr₀₋₁₀Mo_(0-3.5)V_(0-2.5)W_(0-0.15)Ti₀₋₂Nb₂₋₆

In a still preferred embodiment the hardfacing alloy composition isgiven in weight percent by one or a combination of the followingcompositions:

-   Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54)    [alloy 1]-   Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54)    [alloy 2]-   Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂    [alloy3]-   Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.5)Si_(0.59)Ti_(i)V_(0.54)    [alloy 4 ]-   Fe_(ba1)C_(1.2)B₂Mn₁Si_(1.1)Ni_(0.07)Cr_(8.33)Mo_(3.33)V_(0.5)W_(0.07)Ti_(1.83)Nb₄    [alloy 5]-   Fe_(ba1)C₁B_(2.5)Mn₂Si_(1.1)Ni_(0.1)Cr_(8.73)Mo₁V_(0.03)W_(0.03)Ti_(1.91)Nb_(4.47)    [alloy 6]

EXAMPLES Alloys 5 and 6

One of the purposes of designing alloys which possess the non-crackingtraits described within this disclosure is to create a hardfacingmaterial which exhibits very high hardness and wear resistance but isnot prone to re-heat cracking. Two alloys which exhibit both highhardness and resistance to re-heat cracking are alloys 5 and 6. Alloys 5and 6 where produced in the form of welding wires and welded onto astandard 6⅝″ O.D. tool joint in a manner customary to the hardbandprocess used in the oil and gas industry. The feedstock wires were alsomelted into small ingots in an arc-melter, for the purposes of measuringun-diluted hardness and examining microstructure. The results of thehardness measurements for both ingot form and weld bead form are shownin Table 1. Both alloys exhibit high hardness 60 HRC or above, a regionwhich is not typical for crack resistant hardfacing alloys.

TABLE 1 Hardness values of selected disclosed alloys Alloy Form Hardness5 Ingot 63-67 5 Weld Bead 61-63 6 Ingot 59-60

The microstructures of alloy 5 and 6 are shown in FIGS. 7A-B. Bothalloys show a high frequency of carbides within the microstructure whichprovides good hardness and wear resistance, but is typically anindicator for the alloy being prone to cracking. However, both alloyswere deposited via a process typically used in hardbanding as threeconsecutive bands and were free of any cracks. The hardbanding processused reheats existing bead deposits, and is known to generate both dipcracks and circumferential cracks in crack prone alloys of lesserhardness.

1. (canceled)
 2. A hardfacing weld deposit comprising: a hardness of atleast 60 HRC; and a microstructure comprising: an iron-based matrix; andcarbides and/or borides; wherein the carbides and/or borides compriseonly carbides and/or borides which precipitate prior to solidificationof the iron-based matrix.
 3. The deposit of claim 2, wherein thecarbides and/or borides are selected from the group consisting oftitanium boride, niobium carbide, chromium boride, iron-chromium boride,and combinations thereof
 4. The deposit claim 2, wherein the depositdoes not form additional carbides or borides when re-heated to a rangeof 800° C. to 1300° C. for 1 s to 180 s.
 5. The deposit of claim 2,wherein the deposit does not form additional carbides or borides whenre-heated to a range of 900° C. to 1200° C. for 1 s to 180 s.
 6. Thedeposit of claim 2, wherein the deposit does not form additionalcarbides or borides when re-heated to a range of 1000° C. to 1100° C.for 1 s to 180 s.
 7. The deposit of claim 2, wherein the depositcomprises at least one of:Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂;Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.5)Si_(0.59)Ti_(i)V_(0.54);Fe_(ba1)C_(1.2)B₂Mn₁Si_(1.1)Ni_(0.07)Cr_(8.33)Mo_(3.33)V_(0.5)W_(0.07)Ti_(1.83)Nb₄;andFe_(ba1)C₁B_(2.5)Mn₂Si_(1.1)Ni_(0.1)Cr_(8.73)Mo₁V_(0.03)W_(0.03)Ti_(1.91)Nb_(4.47).8. A hardfacing weld deposit comprising: a hardness of at least 60 HRC;and a stable carbide and/or boride structure; wherein a mole fraction ofthe stable carbide and/or boride structure does not change by more than25% when reheated.
 9. The deposit of claim 8, wherein the stable carbideand/or boride structure in the deposit does not change when re-heated toa range of 800° C. to 1300° C. for 1 s to 180 s.
 10. The deposit ofclaim 8, wherein the mole fraction of the stable carbide and/or boridestructure does not change by more than 10% when reheated.
 11. Thedeposit of claim 8, wherein the mole fraction of the stable carbideand/or boride structure does not change by more than 5% when reheated.12. The deposit of claim 8, wherein the deposit further comprises aniron-based matrix, and the deposit possesses a carbide and/or boridethermodynamic stability such that a mole fraction of the carbides and/orborides does not change by more than 25% over a temperature rangebetween room temperature and a solidification temperature of theiron-based matrix.
 13. The deposit of claim 8, wherein the depositfurther comprises an iron-based matrix, and the deposit possesses acarbide and/or boride thermodynamic stability such that any carbidesand/or borides do not form at temperatures above the solidificationtemperature of the iron-based matrix, and are only stable attemperatures below a re-heat temperature range.
 14. The deposit of claim13, wherein the re-heat temperature range is about 800° C. to 1300° C.15. The deposit of claim 13, wherein the re-heat temperature range isabout 900° C. to 1200° C.
 16. The deposit of claim 13, wherein there-heat temperature range is about 1000° C. to 1100° C.
 17. The depositof claim 8, wherein the deposit comprises at least one of:Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂;Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.5)Si_(0.59)Ti_(i)V_(0.54);Fe_(ba1)C_(1.2)B₂Mn₁Si_(1.1)Ni_(0.07)Cr_(8.33)Mo_(3.33)V_(0.5)W_(0.07)Ti_(1.83)Nb₄;andFe_(ba1)C₁B_(2.5)Mn₂Si_(1.1)Ni_(0.1)Cr_(8.73)Mo₁V_(0.03)W_(0.03)Ti_(1.91)Nb_(4.47).18. A hardfacing weld deposit comprising: a hardness of at least 60 HRC;and carbides and/or borides; wherein the carbides and/or boridescomprise an iron concentration of 50 wt. % or less.
 19. The deposit ofclaim 18, wherein the carbides and/or borides are selected from thegroup consisting of niobium carbide, titanium boride, chromium boride,tungsten carbide, molybdenum boride, and vanadium carbide, andcombinations thereof.
 20. A hardfacing weld deposit comprising: ahardness of at least 60 HRC; and an austenite to ferrite transitiontemperature which is outside a re-heat temperature range.
 21. Thedeposit of claim 20, wherein the re-heat temperature range is about 800°C. to 1300° C.
 22. The deposit of claim 20, wherein the re-heattemperature range is about 900° C. to 1200° C.
 23. The deposit of claim20, wherein the re-heat temperature range is about 1000° C. to 1100° C.24. The deposit of claim 20, wherein the deposit comprises at least oneof:Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.54)Si_(0.59)Ti_(0.39)V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V_(0.54);Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb₆Si_(0.59)Ti₁V₂;Fe_(ba1)B_(1.45)C_(0.91)Cr_(4.82)Mn_(1.01)Mo_(3.22)Nb_(4.5)Si_(0.59)Ti_(i)V_(0.54);Fe_(ba1)C_(1.2)B₂Mn₁Si_(1.1)Ni_(0.07)Cr_(8.33)Mo_(3.33)V_(0.5)W_(0.07)Ti_(1.83)Nb₄;andFe_(ba1)C₁B_(2.5)Mn₂Si_(1.1)Ni_(0.1)Cr_(8.73)Mo₁V_(0.03)W_(0.03)Ti_(1.91)Nb_(4.47).